Fiber-reinforced ceramic matrix composite materials are actively being developed for a variety of high-temperature military, aerospace and industrial applications. While possessing high specific strength and toughness, the utility of current ceramic matrix composites are severely limited by their susceptibility to oxidation embrittlement and strength degradation when stressed at or beyond their matrix cracking strength and exposed to high-temperature oxidation. Thus, for the current state of technology, the linear-elastic region represents the “useful” design stress-strain region due to the negative effects caused by environmental degradation of the fiber coating and/or reinforcing fiber at elevated temperatures following the onset of matrix cracking.
Ceramic materials have long been considered potentially beneficial for hot structural component applications in advanced gas turbine and rocket engines, and future high-speed aircraft and atmospheric re-entry vehicles. In general, ceramics have superior high-temperature strength and modulus while having a lower density than metallic materials. The principal disadvantages of ceramics as structural materials are their low failure strain, low fracture toughness and catastrophic brittle failure characteristics. Because of these inherent limitations, monolithic ceramics lack the properties of reliability and durability that are necessary for structural design acceptance. The emerging technology of fiber-reinforced ceramics, or ceramic matrix composites is one promising solution for overcoming the reliability and durability problems associated with monolithic ceramics. By incorporating high-strength, relatively high-modulus fibers into brittle ceramic matrices, combined high strength and high toughness composites can be obtained. Successfully produced ceramic matrix composites exhibit a high degree of non-linear stress-strain behavior with ultimate strengths, failure strains and fracture toughness that are substantially greater than that of the otherwise brittle ceramic matrix.
In order to exploit the benefits of fiber reinforcement in brittle ceramic matrices, it is well recognized that relatively weak fiber/matrix interfacial bond strength is essential for preventing catastrophic failure from propagating matrix cracks. The interface must provide sufficient fiber/matrix bonding for effective load transfer, but must be weak enough to debond and slip in the wake of matrix cracking, leaving the fibers to bridge the cracks and support the far-field applied load. Fiber-reinforced ceramic matrix composites with very high fiber/matrix interfacial bond strengths (usually the result of chemical interaction during manufacture) exhibit brittle failure characteristics similar to that of unreinforced monolithic ceramics by allowing matrix cracks to freely propagate directly through the reinforcing fibers. By reducing the interfacial bond strength, the fiber and matrix are able to debond and slip, thereby promoting the arrest and/or diversion of propagating matrix cracks at/or around the reinforcing fiber. Since crack inhibition/fracture toughness enhancement is the primary advantage of fiber-reinforced ceramic matrix composites, properly engineered fiber coating systems are thus essential for improving the structural performance of these materials. Control of interfacial bonding between the fiber and matrix following manufacture and during service is typically provided by the use of applied fiber coatings.
Fiber-reinforced ceramic matrix composites produced by the chemical vapor infiltration (CVI) process are a particularly promising class of engineered high-temperature structural materials, which are now commercially available. The principal advantage of the CVI process approach for fabricating ceramic composites as compared to other manufacturing methods (e.g., reaction bonding, hot-pressing, melt infiltration, or polymer impregnation/pyrolysis) is the ability to infiltrate and densify geometrically complex, multidirectional fiber preforms to near-net-shape with a ceramic matrix of high purity and controllable stoichiometry without chemically, thermally or mechanically damaging the relatively fragile reinforcing fibers. In addition, because it is a relatively low temperature manufacturing process, high purity refractory matrix materials can be formed (deposited) at a small fraction of their melting temperature (˜Tm/4). Despite the many possible high-temperature ceramic matrix composite systems, however, the number of practical systems is limited by the currently available reinforcing fibers. To date, the majority of high performance ceramic matrix composites produced have primarily been fabricated using carbon and polymer-derived SiC (Nicalon and Hi-Nicalon) fiber reinforcement and CVI-derived SiC matrices.
Carbon fibers offer the highest temperature capability of all current commercially available refractory fibers. Carbon fiber-reinforced SiC ceramics (C/SiC), however, are susceptible to severe strength degradation when exposed to high-temperature oxidizing environments for prolonged periods. This limitation is due to the extensive process-induced matrix microcracking resulting from the relatively large thermal expansion mismatch between the carbon fiber reinforcement and the surrounding SiC matrix. The resultant matrix cracks provide access to environmental intrusion, particularly oxidation, which accelerates the degradation of the compliant fiber coating (e.g., pyrolytic carbon and boron nitride) and the reinforcing fiber. Commercially available small diameter (˜15 μm) ceramic fibers such as Nicalon and Hi-Nicalon microcrystalline SiC, although having limited elevated temperature capability (<1200° C. and <1400° C. respectively) as compared to carbon fiber, exhibit excellent thermomechanical compatibility with SiC matrices. These fibers thus produce composites which are not initially microcracked. Although these ceramic fibers are more oxidation resistant than carbon fibers, the resultant composites also experience irreversible oxidation embrittlement and strength degradation when stressed at or beyond their matrix cracking strength and subsequently exposed to high-temperature oxidation.
Unlike the near-linear tensile stress-strain behavior of the microcracked C/SiC material system, SiC fiber-reinforced SiC matrix composites (SiC/SiC) exhibit highly nonlinear stress-strain characteristics; controlled by the low matrix failure strain relative to the reinforcing fiber. As the composite is loaded in tension, it deforms linear-elastically up to the onset of matrix cracking. The tension threshold at which the onset of matrix cracking occurs designates the “proportional limit” of the material. That is, when the applied tensile strain reaches the failure strain of the unreinforced matrix, ideally assuming negligible residual thermoelastic effects, transverse matrix cracks initiate and propagate rapidly across the composite, leaving the fibers to bridge the cracks while supporting the far-field applied load. Continued loading beyond the onset of matrix cracking results in the formation of many regularly spaced matrix cracks (i.e., multiple matrix cracking) typically accompanied by a nonlinear decrease in composite stiffness. This strain-induced compliance behavior is the result of the diminishing contribution of the matrix modulus with increased multiple matrix cracking and fiber/matrix debonding. The diminishing stiffness behavior becomes more significant with increasing applied strain to the point where the elastic modulus of the composite is primarily dominated by the reinforcing fibers. As the composite strain approaches the fiber failure strain, the fibers progressively fracture, designating the ultimate strength of the composite. For most practical structural applications, however, the linear-elastic region represents the “useful” design stress-strain region due to the negative effects of hysteresis and environmental degradation of the fiber coating and/or fiber reinforcement at elevated temperatures occurring after matrix microcracking. Matrix microcracking is therefore a fundamental life-limiting issue for ceramic matrix composites being considered for use in extended-life thermostructural applications.
From an engineering mechanics standpoint, the high elastic modulus of the CVI-derived SiC matrix relative to the reinforcing fiber is a disadvantage for load transfer. For a Nicalon SiC fiber-reinforced/CVI SiC matrix composite with a 40 volume-percent fiber loading, nominally 70% of the applied load is carried by the matrix prior to the onset of matrix cracking. For an equal volumetric loading of higher modulus Hi-Nicalon SiC fiber, about 60% of the applied load is initially carried by the matrix. The relatively high matrix stiffness is thus a disadvantage from the standpoint of matrix microcracking. The tensile strain at which the onset of matrix cracking occurs in both Nicalon and Hi-Nicalon reinforced SiC matrix composites is typically ˜0.04%, and rarely exceeds 0.05%, as this is an intrinsic property of the CVI-derived brittle matrix. The corresponding tensile matrix cracking strengths typically range between 60 and 80 MPa, respectively. In turn, the useful design strengths for both Nicalon and Hi-Nicalon SiC fiber-reinforced/CVI SiC composites are only about 30% of their respective fiber-dominated ultimate strengths. Thus, the early onset of matrix microcracking and subsequent oxidation embrittlement and strength degradation is a primary performance limitation of current state-of-the-art materials. Fundamentally, it would be desirable for a fiber-reinforced ceramic composite material to have a useful design strength significantly greater than 30% of its ultimate strength.
Although the oxidation embrittlement problem in fiber-reinforced ceramic matrix composite materials may eventually be controlled via advanced fiber coating and/or matrix oxidation inhibition approaches, it will nevertheless still be desirable to increase the elastic limit of the composite to reduce potential fatigue, hysteresis and other complex nonlinear material behavioral effects. Once the composite elastic limit is exceeded, the structural designer is faced with using a nonlinear and potentially time-dependent microcracked material system. One, however, nontrivial approach towards increasing the matrix cracking strength in ceramic matrix composites is by increasing the mechanical properties (e.g., strength and fracture toughness) of the matrix constituent. This approach has been successfully demonstrated by the current inventor via microstructural engineering of the SiC matrix into a strong/tough nanolayered composite constituent.
A nanolayered composite comprises a compositionally modulated microlaminate consisting of periodically alternating layers (lamellae) of two or more material constituents. The thickness of each layer range from about one molecular monolayer (˜1 nm) to a thickness approaching the upper limit of very fine grain refinement achievable from current state-of-the-art materials processing techniques (>150 nm). These materials can be engineered to exhibit remarkable mechanical, tribological, thermal, and/or electrical properties that are uniquely different from those of the individual constituents. In particular, strength can be enhanced over currently available courser grained materials by an order of magnitude or more. Also of importance is that a conceivably wide range of refractory metal and ceramic materials can be engineered into such nanostructural composites suitable for extreme environmental structural applications.
Early efforts (over two decades ago) by researchers at the Chemetal/San Fernando Laboratories (SFL) led to the discovery of a unique form of chemical vapor deposited SiC. While attempting to deposit “massive” bodies of SiC, unanticipated thermochemical process instabilities (i.e., chugging) within the “cold-wall” chemical vapor deposition (CVD) reactor resulted in producing a material with an unusual layered microstructure. This material was found to be composed of alternating lamellae of SiC and elemental silicon (Si), ranging in thickness from 10 to 20 nm and 1 to 2 nm, respectively. Reported properties for this SiC/Si material included flexural strengths, elastic moduli, fracture toughness, and hardness which exceeded 4000 MPa, 450 GPa, 6-12 MPa√m, and 45 GPa, respectively. Numerous subsequent evaluations by government laboratory scientists Dutta, Graham, Rice, and Mendiratta confirmed these astonishing results. Dutta, S., R. Rice, H. Graham, and M. Mendiratta, Characterization and Properties of Controlled Nucleation Thermochemical Deposited (CNTD) Silicon Carbide, NASA TECH. MEMO. 79277, presented at the 80th Annual Meeting of the American Ceramic Society. This work resulted in the issuance of a number of domestic and foreign patents for which the process was coined “Controlled Nucleation Thermochemical Deposition”, or CNTD.
Despite the extraordinary mechanical and physical properties of this termed “ultra-structured” material, however, commercialization was hindered by problems of reproducibility. In short, processing difficulties associated with the inability to control the naturally occurring chemical instability within the cold-wall reactor during deposition prevented this product from becoming commercially successful. Specifically, the uncontrollable, and not well-understood “cyclic” instability phenomenon was not easily scaled to larger or hot-wall reactors, resulting in low yield, poor reproducibility and poor process economics. Accordingly, it would be desirable to be able to artificially reproduce the beneficial effects of this CNTD process such that it could be effectively used in hot-wall CVD reactors (which can accommodate large batch quantities of dissimilar parts) by relatively simple and controllable mechanical means.
It is also known that layered fiber coatings can be applied to fibrous preforms in such a way to increase the oxidation resistance of the resultant ceramic matrix composite, while preserving desirable strength and toughness. Particularly, by controlling the flow of the chemical precursors (i.e., chemically reactant precursor gases or gasified liquids) during application of a ceramic coating onto a preform of refractory fibers, wherein two or more independent chemical precursors are periodically turned on and off at prescribed intervals, it has been shown that the resultant microlayered coating produced creates an inherently oxidation-resistant fiber coating material. FIG. 1 shows a backscattered electron image (BEI) of an advanced “multilayered SiC fiber coating” developed by the current inventor and applied using the technique of cyclic “throttling” of the chemical precursors during the deposition of the fiber coating. Details of this multilayered fiber coating and processing method is more fully described in U.S. Pat. Nos. 5,455,106 and 5,545,435, the disclosures of which are incorporated herein by reference. The microlayered fiber coating system (deposited on a ˜15 μm Hi-Nicalon SiC fiber) shown in FIG. 1 was engineered to mitigate the inherent problems of oxidation resistance plaguing currently available PyC and BN fiber coatings for structural ceramic matrix composites. This was achieved by successfully tailoring the desired mechanical characteristics (e.g., interfacial shear strength and compliance) of the multilayered SiC coating system via microstructural engineering necessary to enhanced strength and toughness of the resultant composite.
These patent disclosures, however, did not address the problem of increasing the strength and/or toughness of the matrix constituent itself and were directed instead to the product and method of depositing the oxidation-resistant multilayered ceramic fiber coating material. These patent disclosures describe depositing microlayers having a primary layer thickness of between 500 and 5000 nanometers, which is considered too thick to increase the inherent strength of the material produced as will be described later. Accordingly, it would be desirable to increase the strength and toughness of the ceramic matrix constituent in order to enhance the matrix cracking resistance in the resulting composite.
To better understand the unique behavior of CVI/CVD nanostructures, it is useful to briefly describe the morphology of these engineered materials. Nanolayered composites are produced by depositing a layer of the primary, or major constituent species with a thickness of on the order of a few tens of nanometers (10-100 nm), followed by a layer of minor species with a thickness of about an order of magnitude less (1-10 nm). Deposition durations for each layer are very short, ranging from a few seconds to a few tenths of a second. The exact deposition durations are dependent on the deposition rates of the respective major and minor species derived from a given process. The process is then cyclically repeated until the desired thickness of the body is achieved. Although there may appear to be great flexibility in the selection of the secondary nanolayering constituent(s), they must be carefully selected based on their (1) known ability to effectively interrupt the deposition epitaxy of the major constituent, thereby increasing strength and thermal shock resistance by controlling grain refinement; and (2) propensity to provide beneficial elastic modulus mismatch, thereby further increasing fracture toughness by limiting dislocation motion.
In conventional CVI/CVD-deposited materials, high-purity crystallites nucleate on the heated substrate surface (e.g., part to be coated or fibers) and then grow epitaxially in a direction perpendicular (i.e., normal) to the heated substrate; most often through the entire thickness of the deposit. The crystallites thus coarsen and weaken with increasing thickness during the growth process. In the current invention (e.g., CVI/CVD nanolayering process), the major disadvantages of conventional CVI/CVD are eliminated. The first crystallites nucleate and start to grow, competing for a preferred minimum energy orientation. Before they ever have a chance to become oriented, the growth is interrupted by the deposition of the second material. This secondary layer is deposited so thinly that its crystallites do not have a chance to grow, and thus do not achieve any preferred orientation. When the cycle is repeated, the crystallites of the primary material must re-nucleate and the process of nucleation/interruption is repeated; thus, the grains in the deposited material never have a chance to coarsen. When engineered successfully, this process has been shown to result in producing materials with significantly increased strength and hardness; beyond that predicted by the teachings of Hall-Petch. Crystals of about 5 μm are considered fine by most materials scientists and engineers; and those of 1 μm, extremely fine. Virtually no structural components have ever been produced by conventional methods with grains less than 0.4 μm (400 nm) in size. The CVI/CVD nanolayering process developed in the current invention provides the ability to produce highly uniform microstructures with grain sizes of between 1 and 100 nm. Although engineering properties are improved by grain refinement in general, it is not until the crystallites are maintained to less than ˜150 nm that dramatic improvements to near-theoretical strengths are observed.
FIG. 2 shows a microstructural example of a nanolayered CVD SiC material developed by the current inventor. A tensile strength enhancement of nearly one (1) order of magnitude (˜8×) over conventional CVD SiC has been experimentally demonstrated.